PREPARATION OF: I. INTERCALATIVE METAL OXIDE/CONDUCTIVE POLYMER COMPOSITES AS ELECTRODE MATERIALS FOR RECHARGEABLE BATTERIES; II. SODIUM RICH MANGANESE OXIDE HYDRATE WITH CAPACITY FOR AQUEOUS Na-ion ELECTROCHEMICAL ENERGY STORAGE

ABSTRACT

The present invention is directed at intercalative metal oxide/conductive polymer composites suitable for use as electrode materials for rechargeable batteries. The composites can be prepared by agitation of the metal oxide and the conductive polymer in aqueous media. The present invention is also directed at a sodium rich layered manganese oxide hydrate prepared by annealing manganese (II, III) oxide and sodium hydroxide. The sodium rich manganese (III, IV) oxide so formed indicates an enhanced capacity for Na-ion storage suitable for the use of electrode materials for aqueous energy storage.

CROSS-REFERENCE TO RELATED APPLICATIONS

The present application is a by-pass continuation of PCT/US18/60677 filed on Nov. 13, 2018 and claims the benefit of the filing date of U.S. Provisional Application, Ser. No. 62/585,031, filed Nov. 13, 2017, both of which are fully incorporated herein by reference.

GOVERNMENT RIGHTS CLAUSE

This invention was made with government support under Prime Contract No. DE-SC0010286 awarded by the U.S. Department of Energy. The government has certain rights in the invention.

FIELD

The present invention is directed at intercalative metal oxide/conductive polymer composites suitable for use as electrode materials for rechargeable batteries. The composites can be prepared by agitation of the metal oxide and the conductive polymer in aqueous media. The present invention is also directed at a sodium rich layered manganese oxide hydrate prepared by annealing manganese (II, III) oxide and sodium hydroxide. The sodium rich manganese (III, IV) oxide so formed indicates an enhanced capacity for Na-ion storage suitable for the use of electrode materials for aqueous energy storage.

BACKGROUND

The preparation of effective electrode materials especially cathode materials for rechargeable energy storage devices including supercapacitors and batteries have attracted increased attention. Many metal oxide, for example vanadium pentoxide, manganese oxide, cobalt oxide are the most promising layered electrode materials for rechargeable lithium/sodium/potassium batteries because of its high-energy storage capacities. However, the capacity retention of these electrode materials upon cycling and power performance (time required to charge and discharge the devices) are unsatisfactory, partially due to their poor electrical conductivity. Several approaches have been developed on the preparation of new electrode materials to ameliorate these limitations, including simply mixing electrode material with carbon black or coating the electrode material with more conductive layers. In particular, the preparation of metal oxide/conductive polymer composite is a very rapidly developing area of electrode materials development. Compared with simple mixtures of metal oxide and conducting polymer, preparation of intercalative metal oxide/conducting polymers is a very promising approach to making electrode materials, where conductive polymer is inserted between metal oxide layers, so that semiconductive metal oxide materials are in close proximity to conductive polymer. However, typical preparation routines for intercalative composite materials include in situ polymerization, or require assistance of heat or even microwave radiation, and therefore involve relatively long processing time, arduous preparation procedures and have difficulty in scale-up production. It is, therefore, necessary to develop simpler and more energy-efficient methods of synthesis method in this regard.

In related context, developing electrochemical energy storage (EES) technologies using safe and earth-abundant materials becomes increasingly attractive for economically storing the electric power generated by solar and wind. Aqueous EES devices using Na-ions as charge carriers have been promising alternatives to non-aqueous lithium-ion batteries (LIBs) for its low cost, high safety and availability of Na sources in terrestrial reserves. However, the storage of Na-ions is challenging because of its relatively large ionic radius, so that LIB host materials (especially cathode) usually with a close-packed array of oxide ion may not be able to accommodate the Na-ion for reversible insertion and extraction.

Two design principles have been used to tackle the intercalation of Na-ion. One is the replacement of oxygen anions (O²⁻) with anions that have weaker bonding with metal cations, so that cations are sufficiently mobile in the electrode. Recent studies show promise of hexacyano ions (C≡N)₆ ⁶⁻ based electrode materials for Na- and K-ions storage due to their weakened bonding between cyanide (C≡N)⁻ and cations. Some reports indicate that potassium copper hexacyanoferrate and its analogues can function as stable electrode materials for aqueous K- and Na-ion storage. Sodium manganese hexacyanoferrate has been reported by to show relatively good energy performance and cycling life in a non-aqueous electrolyte.

Another approach to design a Na-ion electrode is to introduce a relatively large interstitial host framework. These materials with planar or zigzag layers show different polymorphs (P2, P3 or O3 symmetry) with respect to the sites of the intercalated alkali ions by altering the stacking of transition metal-oxygen octahedra ([MO₆]). However, the mechanistic understanding of storage of Na-ion inside various host materials is still far from settled. This is largely due to the different intercalation chemistry of Na-ions from that of Li-ions so that the fundamental understanding obtained from Li-ion storage may not be directly applied to Na-ion.

Recently, efforts have been devoted to the studies of the storage mechanisms of various alkali ions. For example, O3-type layered LiMnO₂ (ABC oxygen stacking) suffered from degradation to spinel structure and thus the impaired capacity for Li-storage due to the migration of Mn ions during the cycling. In contrast, the NaMnO₂ counterpart had a high energy barrier for Mn ion migration, which prevented cation mixing and thus sustained the layered structure during the Na-ion intercalation/deintercalation. Compared with O3-type NaMnO₂, birnessite δ-MnO₂ also has a layered structure containing two-dimensional sheets of edge-shared MnO₆ octahedra with a general formula of A_(x)MnO₂.H₂O (A: H⁺, Li⁺, Na⁺, or K⁺; X: usually less than 0.2).¹⁻⁵ The studies of birnessite electrode in aqueous electrolyte have been reported. However, even though the birnessite has a rather large interlayer distance (˜7 Å), the storage capacity was typically low for aqueous Na-ion storage (<60 mA h g⁻¹) due to the limited potential window (˜1.2 V) and ineffective redox process. Much less work on improving the aqueous Na-ion storage capacity in birnessite has been reported to date.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 shows photos of bulk V₂O₅ in water, simple mixture of V₂O₅, PEDOT:PSS and water, as well as intercalative V₂O₅/PEDOT:PSS composite. The color change reflects structural differences of intercalative V₂O₅/PEDOT:PSS composite compared with V₂O₅ or simple mixture of V₂O₅ and PEDOT:PSS.

FIG. 2 shows XRD patterns of commercial V₂O₅ (the precursor for making intercalative V₂O₅/PEDOT:PSS composite) and intercalative V₂O₅/PEDOT:PSS composite.

FIG. 3 shows SEM and TEM images of commercial V₂O₅ (the precursor for making intercalative V₂O₅/PEDOT:PSS composite) and intercalative V₂O₅/PEDOT:PSS composite. The latter shows distinct morphological changes from bulk particles to layered structure.

FIG. 4 shows neutron pair distribution function (PDF) analysis of commercial V₂O₅ (the precursor for making intercalative V₂O₅/PEDOT:PSS composite), and intercalative V₂O₅/PEDOT:PSS composite. The latter shows distinctly decreased coherent length, corresponding well to the formation of layered structure.

FIG. 5A shows electrochemical tests of intercalative V₂O₅/PEDOT:PSS composite at the scan rate ranging from 10 mV/sec to 1000 mV/sec in an aqueous KCI electrolyte.

FIG. 5B shows electrochemical tests of bulk V₂O₅, simple mixture of V₂O₅ and PEDOT:PSS, as well as intercalative V₂O₅/PEDOT:PSS composite. The latter showed enhanced capacity (more than three times).

FIG. 6A shows TEM images of LiMnO₂ bulk material.

FIG. 6B shows intercalative LiMnO₂/PEDOT:PSS nanosheet.

FIG. 7 shows cyclic voltammetry measurements of bulk LiMnO₂ materials.

FIG. 8 shows cyclic voltammetry measurements of intercalative LiMnO₂/PEDOT:PSS nanosheet.

FIG. 9 shows a comparison of specific capacity toward K-ion electrochemical storage of intercalative LiMnO₂/PEDOT:PSS nanosheets and bulk LiMnO₂ materials where intercalative nanosheets show at least three times higher capacity.

FIG. 10 shows EDS characterizations of Na₆MnO_(x) materials with (a) Mn₅O₈ obtained by directly heating Mn₃O₄ nanoparticles, (b)N_(0.13)MnO_(1.74)-H₂O, (c) Na_(0.25)MnO_(1.84)-H₂O, and (d) Na_(0.29)MnO₂—H₂O obtained by the thermal solid-state reactions of NaOH and Mn₃O₄ as the molar ratios of 0.5:1, 1:1, 2:1, respectively.

FIG. 11 shows (a) EDS characterization of Na_((δ))MnO_(x)—H₂O (Na_(0.13)MnO_(1.74)-H₂O; Na_(0.25)MnO_(1.84)-H₂O and Na_(0.29)MnO₂—H₂O) materials obtained by the thermal solid-state reaction of NaOH and Mn₃O₄ as a molar ratio of 4:1; (b) the summaried atomic ratios of sodium to manganese for various Na₆MnO_(x)—H₂O materials; and (c) XRD patterns of Na_(0.29)MnO₂—H₂O materials syntheized via solid state reaction of NaOH:Mn₃O₄ as molar ratios of 2:1 and 4:1, indicating the forming of a stable Na_(0.29)MnO₂ material even with an increased amount of NaOH precursor.

FIG. 12 shows EDS characterization of Na_(δ)MnO₂—H₂O birnessite synthesized via the wet chemistry method showing the Na/Mn atomic is 0.17:1 with the comparision to that of other works in Table 1.

FIG. 13 shows (a) TEM image of of Na_(0.29)MnO₂.H₂O materials; (b) Experimental neutron PDFs of Na₆MnO_(x)—H₂O materials, where coherent lengths of the materials decreased as the Na concentration increased;(c) Phase percentage of Na_(0.29)MnO₂—H₂O in Na_(δ)MnO_(x).H₂O materials obtained from neutron PDF analysis.

FIG. 14 shows TEM characterizations of Na_(δ)MnO_(x)—H₂O materials with (a) Mn₅O₈, (b)Na_(0.13)MnO_(1.74)-H₂O, (c) Na_(0.25)MnO_(1.84)-H₂O.

FIG. 15 shows neutron PDF analysis of (a) Mn₅O₈, (b) Na_(0.13)MnO_(1.74).H₂O, (c) Na_(0.25)MnO_(1.84).H₂O and (d) Na_(0.29)MnO₂—H₂O materials.

FIG. 16 shows XRD patterns of Na_(δ)MnO_(x) materials compared with Mn₅O₈ and MnO₂ standards indicating the structural evolution from Mn₅O₈ to MnO₂ as sodium concentrations increased, and the Rietveld refinement analyese of XRD pattern of Na_(0.29)MnO₂—H₂O.

FIG. 17 shows neutron PDFs of Na_(δ)MnO_(x) materials normalized by the intensity of the peak at 1.9 Å associated with Mn—O pair. The atom pair associated with each peak (P1 to P7) can be attributed to (a) water, (b, d, e, g, i) Mn₅O₈ polyhedra in black and (c, f, h) MnO₂ polyhedra in blue.

FIG. 18 shows the proposed formation mechanism for Na_(0.29)MnO₂—H₂O driven by solid-state Na-ion intercalation in (Mn²⁺: green; Mn³⁺: orange; Mn⁴⁺: purple; Na⁺: brown; O: red).

FIG. 19 shows XRD characterizations of Na_(0.29)MnO_(x) materials with the thermal solid-state reacted Na_(0.29)MnO₂ and washed Na_(0.29)MnO₂ (the diffraction peaks from Mn₃O₄ are indexed in black and the black solid dot from NaOH, the resulting heated anhydrous Na_(0.29)MnO₂ and washed Na_(0.29)MnO₂—H₂O are labelled as well).

FIG. 20 shows electrochemical CV measurements in half-cells. Cyclic Voltammetry (CV) scans of (a) Mn₅O₈, (b) Na_(0.13)MnO_(1.74)-H₂O, (c) Na_(0.25)MnO_(1.84)-H₂O, and (d) Na_(0.29)MnO₂—H₂O between −1.25 V to 1.25 V (vs Ag/AgCl) in 0.1 M Na₂SO₄ electrolyte at the scan rates from 5 to 1000 mV s⁻¹.

FIG. 21 shows electrochemical half-cell measurements with (a) Cyclic Voltammetry (CV) scans of Na_(0.29)MnO₂.H₂O material between −1.25 V to 1.25 V (vs Ag/AgCl) at various scan rates in 0.1 M Na₂SO₄ electrolyte; (b) calculated specific charge storage capacities of sodium-manganese oxides as a function of scan rates; Symmetric full-cell measurements with (c) discharge electrode capacities of Na_(0.29)MnO₂.H₂O material at the various current densities of 1 A g⁻¹, 2 A g⁻¹, 5 A g⁻¹ and 10 A g⁻¹ (after 5000 galvonstatic charge and discharge process unless specified otherwise); (d) electrode capacites of Na_(0.29)MnO₂—H₂O as a function of cycle number up to 5000 at the current densities from 1 A g⁻¹ to 10 A g⁻¹; (e) coulombic and energy efficiencies of Na_(0.29)MnO₂—H₂O at various current densities as a function of current density (1, 2, 5 and 10 A g⁻¹); (f) Ragone plot with gravimetric specific energy and power of the symmetric Na_(0.29)MnO₂.H₂O fullcell after 5000 galvanostatic cycles. The aqueous (empty symbols) and nonaqueous (solid symbols) devices are reported, and the gravimetric specific energy and power are calculated by the mass of electrode materials except the Panasonic (17500) Li-ion batteries.

FIG. 22 shows CVs of sodium-manganese oxides at the scan rate of 50 mV s⁻¹, showing the anodic peak of Na_(0.29)MnO₂—H₂O shifted to a lower potential and the cathodic peak shifed to a higher potential compared with those of other materials as the Na concentration increased.

FIG. 23 shows diffusivity measurements of Na_(δ)MnO_(x) materials with (a) Mn₅O₈, (b) Na_(0.13)MnO_(1.74)-H2O, (c) Na_(0.25)MnO_(1.84)-H₂O, (d) Na_(0.29)MnO₂—H₂O, (e) the relaxation steps of Na_(δ)MnO_(x) materials and (f) (E₀-E) as a function of (1/t^(1/2)) curves for the slope calculations, where E₀ is the open circuit voltage.

FIG. 24 shows electrode capacities of Na_(0.29)MnO₂.H₂O as a function of voltage at the current denties of (a) 1 A g⁻¹, (b) 2 A g⁻¹, (c) 5 A g⁻¹ and (d)10 A g⁻¹ during galvanostatic charge and discharge process.

FIG. 25 shows (a) CV scans of disorder Na_(0.29)MnO₂—H₂O, high temperature treated Na_(0.29)MnO₂—H₂O and commercial MnO₂ bulk materials at the scan rate of 5 mV s⁻¹ in a 2.5 V potential window in half-cell; (b) calculated Tafel slopes at the scan rate of 5 mV s⁻¹; (c) Specific capacities at the scan rate of 5, 10 and 50 mV s⁻¹; (d) Oxygen K-edge sXAS of electrochemically cycled Na_(0.29)MnO₂, Mn₅O₈ and anhydrous commercial MnO₂ bulk materials.

FIG. 26 shows CV scans of (a) disordered Na_(0.29)MnO₂.H₂O, (b) high temperature treated Na_(0.29)MnO₂—H₂O and (c) commercial MnO₂ bulk materials between −1.25 V to 1.25 V (vs Ag/AgCl) in 0.1 M Na₂SO₄ electrolyte at the scan rates of 5, 10 and 50 mV s⁻¹; Tafel plots of (d, e, f) HER and (g, h, i) OER for disordered Na_(0.29)MnO₂, high temperature treated Na_(0.29)MnO₂—H₂O and commercial MnO₂ bulk at scan rates of 5, 10 and 50 mV s⁻¹, and the summaried Tafel slopes of (j) HER and (k) OER.

DETAILED DESCRIPTION

The present invention is directed at preparation of intercalative (layered) metal oxide/conductive polymer composites as electrode materials for rechargeable batteries. Preferred metal oxides include those oxides that can form a two-dimensional plane with relatively strong in-plane chemical bonding with a dissociation energy of 4 to 7 electron volts (eV) (the average vanadium-oxygen ionic bonding is around 6.7 eV and manganese-oxygen ionic bonding is around 4.2 eV) and relatively weak Van der Waals bonding between planes with a dissociation energy of about 0.01 eV. Preferred examples include V₂O₅ and LiMnO₂. It is contemplated that other suitable oxides may include TiO₂, MoO₂, MoO₃, Nb₂O₅ and LiCoO₂. The conductive polymer herein is preferably a positively charged polymeric ionomer in combination with a negatively charged polymeric ionomer. Reference to ionomer herein is to be understood as a charged polymer stabilized by ionic cross-links.

One particularly preferred conductive polymer includes as the positive charged ionomer poly(3,4-ethylene dioxythiophene) (PEDOT) in combination with the negatively charged ionomer poly(styrenesulfonate) (PSS). The conductive polymer may therefore be conveniently identified as PEDOT:PSS. In PEDOT:PSS, part of the sulfonyl groups are deprotonated and carry a negative charge. The PEDOT is a conjugated polymer and carries positive charges based upon polythiophene.

The intercalative structure (layering of the oxide and conductive polymer) is preferably achieved herein by agitation (e.g. stirring) of the metal oxide and the conductive polymer (i.e. positively charged polymeric ionomer in combination with negatively charged polymeric ionomer) in aqueous media. For example, the intercalative structure herein was observed to form when agitating the metal oxide with the PEDOT:PSS in water for an extended period of time, such as for 100 to 200 hours. Preferably, the metal oxide to conductive polymer weight ratio is in the range of 1:1 to 8:1, more preferably 3:1 to 5:1. One particularly preferred weight ratio of metal oxide to conductive polymer is 4:1.

The above procedure is a relatively scalable synthetic routine and is preferably carried on at room temperature without heat or radiation. The resulting nanocomposites have been characterized by powder X-ray diffraction, Raman spectroscopy and transmission electronic microscope analyses. The thickness of the layered structure is contemplated to fall in the range of 1-5 nm up to 30 nm.

Furthermore the application potential of the nanocomposites herein have been tested in an aqueous sodium batteries test, which display some synergistic effects between the metal oxides (V₂O₅, NaMnO₂) and the intercalative conductive polymer (PEDOT:PSS). The results showed that intercalative metal oxide/conductive polymer composites show 100% to 400% enhanced capacity, as well as much improved power performance.

The preparation of intercalative V2O5/PEDOT:PSS nanocomposite was as noted preferably conducted in aqueous solution at room temperature using the conductive polymer poly(3,4-ethylenedioxythiophene)-poly(styrenesulfonate) [PEDOT:PSS]. A mixture of 100 mg of commercial V2O5 bulk material and 25 mg of PEDOT:PSS were submerged in 6 mL of deionized (DI) water in a scintillation vial. The mixture was vigorously stirred at approximately 500 rpm for one week, accompanied with a noticeable color change (FIG. 1). The dispersion of the commercially available V₂O₅ powder in DI water is light yellow in color. The conductive polymer PEDOT:PSS, a light blue color, is added to the V₂O₅ giving the mixture a slightly green tinge. Over the one week mixing period, a color change from the green-yellow color of the mixture to a dark green color occurs. The oxidation state of the V⁵⁺ likely reduced to a V^(4+/5+) mixture due to the interaction with PEDOT:PSS.

Through XRD a significant change can be clearly seen that the crystal structure and morphology is affected by the interaction with the PEDOT:PSS (FIG. 2). The V₂O₅ is initially highly crystalline having relatively strong Bragg features, orthorhombic cell with Pmmn symmetry. However after mixing in DI water with the PEDOT:PSS for two weeks a layered structure is formed. The PEDOT:PSS appears to insert into the interlayer gap of the bulk V₂O₅ materials having a weak Van der Waals force between layers, which finally facilitates the formation of the intercalative V₂O₅/PEDOT:PSS nanosheets through a exfoliation process. The diffraction pattern of the exfoliated layered structure produces only a few small Bragg peaks in the XRD pattern. The d-spacing of the layered peak d₀₀₁=13.67 Å (3.05° 2θ).

SEM and TEM images show that the V₂O₅ before exfoliation appears as particles and after exfoliation the V₂O5-PEDOT:PSS appear as ribbon like strands appearing to pull away from each other and the bulk particles (FIG. 3).

Neutron pair distribution function (PDF) analysis was chosen to elucidate the interaction between the V₂O₅ and PEDOT:PSS due to the lack of Bragg features in the diffraction pattern (FIG. 4). The exfoliated layer phase was identified as a bi-layered V₂O₅ structure, having a coherence length of 15 Å, a significant decrease as compare to the bulk phase. There was still a detectable bulk phase in the PDF, 8% by mass. A simplified model currently has atoms with the correct ratio of elements for the PEDOT:PSS in the interlayer spacing but not in the correct conformation. The low r-space data is dominated by water and the PEDOT:PSS. The first strong negative peak at ˜1.01 Å is due to both the O-H interaction water (0.96 Å) and the interaction of the polymer C—H (1.1 Å). The strong positive peak at ˜1.45 Å is the C—C interaction from the conductive polymer and a small portion is due to the H—H correlation from water. The second negative peak at 2.09 Å is C—H interaction but a carbon atom interacting with a hydrogen atom bonded to a nearest neighbor carbon atom.

The intercalative V₂O₅-PEDOT:PSS nanocomposite showed an increase in capacitance as compared to a mixture of the V₂O₅ and PEDOT:PSS (FIG. 5B). The capacity of 160 mAh/g at a scan rate of 10 mV/s is among the best reported K-ion electrochemical storage using aqueous electrolytes. As shown in FIG. 5B, over the scan rate of 10 mV/s the intercalative V₂O₅-PEDOT:PSS nanocomposite indicates a capacity of greater than 75 mAh/g, or in the range of 75 MAh/g to 160 mAh/g. The redox features remain visible though all scans rates tested from 10 1000 mV/s, some shoulders start to appear at higher scan rates. FIGS. 6-9 show the corresponding evaluations and comparisons of bulk LiMnO₂ and intercalative LiMnO₂/PEDOT:PSS nanosheets material prepared by the procedure herein of agitation of LiMnO₂ in water in the present of PEDOT:PSS. Once again, as shown in FIG. 9, the intercalative LiMnO₂/PEDOT:PSS nanosheets indicate at least three times higher electrochemical storage capacity as compared to bulk and non-layered LiMnO₂. More specifically, as show in FIG. 9, the intercalative LiMnO₂/PEDOT:PSS indicates the following: (1) at a scan rate of 10-20 mV/sec, a specific capacity of greater than or equal to 60 mAh/g or in the range of 60 mAh/g to 70 mAh/g; (2) at a scan rate of 10-500 mV/sec, a specific capacity in the range of 20 mAh/g to 70 mAh/g.

Turning next to the sodium-rich manganese oxide hydrates with capacity for aqueous sodium ion electrochemical energy storage, the following is noted. Preferably, Na-rich MnO₂—H₂O suitable for use for aqueous Na-ion storage can be made in the solid state, preferably by annealing the mixture of Mn₃O₄ and NaOH, involving conversion from Mn₃O₄ spinel to an ordered Mn₅O₈ layered structure and finally to Na-rich MnO₂—H₂O driven by Na-ion insertion. The Na-rich manganese oxide hydrate herein is represented by the formula Na_((δ))MnO_(x)—H₂O wherein δ has a value greater than 0.17, or more preferably, in the range of >0.17 to 0.29; and x has a value in the range of 1.74 to 2.0. The reaction was confirmed by neutron total scattering measurements and pair distribution function (PDF) analysis. Storage capacity up to about 150 mA h g⁻¹ is observed through increase of the potential window and promotion of the redox charge transfer process towards the aqueous Na-ion storage. It should be noted that the Mn₃O₄ precursor is a manganese (II, III) oxide, where the valences of the Mn element are 2+ and 3+. In addition, the Na_((δ))MnO_(x)—H₂O is a manganese (III, IV) oxide which includes Mn⁴⁺ and Mn³⁺.

The resulting Na_(0.29)MnO₂—H₂O material exhibits a relatively high overpotential (˜0.6 V) towards oxygen and hydrogen evolution reactions and therefor enables a kinetically stable potential window of 2.5 V in the half-cell in an aqueous electrolyte without gas evolution. Moreover, the Na-rich structure improves diffusion-limited redox charge storage encouraging up to a 0.41 electron transfer reaction. Overall, the resulting Na_(0.29)MnO₂—H₂O demonstrates a reversible capacity of about 130 to 160 mA h g⁻¹ (a scan rate of 5 mV s⁻¹ in the half-cell) in aqueous Na-ion storage, a high energy density of 20 to 30 Wh kg ⁻¹ (a rate of 23° C. in a full-cell), and a relatively good cycling life (70 to 100 mAh g⁻¹ after 5000 cycles at an electric current rate of 1 A g⁻¹ in a full-cell).

Different from wet chemistry methods involving the oxidation of Mn²⁺ or reduction of permanganate at room temperature, Na_((δ))MnO_(x)—H₂O materials (Na_(0.13)MnO_(1.74)-H₂O; Na_(0.25)MnO_(1.84)-H₂O and Na_(0.29)MnO₂—H₂O) were preferably prepared at 270° C. in the air via a solid-state reaction between NaOH and Mn₃O₄nanoparticles, followed by water rinsing. The temperature range for the solid state reaction may fall in the range of 200° C. to 400° C., more preferably 250° C. to 300° C. The Mn₃O₄ nanoparticles may have a size range of 10 to 30 nm. By altering the molar ratios between NaOH and Mn₃O₄ from 0 to 2, various sodium manganese oxides (Na_(0.13)MnO_(1.74)-H₂O; Na_(0.25)MnO_(1.84)-H₂O and Na_(0.29)MnO₂—H₂O) were prepared, verified by energy dispersive X-ray spectroscopy (EDS) measurement (FIG. 10). Atomic ratios of Na/Mn remained the maximum value of 0.29 when the ratio of NaOH: Mn₃O₄ increased to 4:1 (FIG. 11). Such a high Na/Mn ratio of 0.29 can only be achieved via the solid-state annealing, while birnessite made via a wet chemistry approach has a Na/Mn ratio of 0.17 (FIG. 12, Table 1). The morphology of Na_((δ))MnO_(x)—H₂O materials evolved from faceted nanoparticles, to a mixture of layers and particles, and finally to a complete layered structure with a planar dimension up to 200 nm, when the Na concentration (δ) increased from 0 to 0.29 (FIGS. 22a & 23).

FIG. 13b shows the PDF data for various Na_((δ))MnO_(x)—H₂O materials (Na_(0.13)MnO_(1.74).-H₂O; Na_(0.25)MnO_(1.84)-H₂O and Na_(0.29)MnO₂—H₂O) obtained from neutron total scattering measurements. The lattice parameters obtained after refinement were shown in FIG. 15 and Tables 2-5. Unlike Rietveld refinement that only analyzes the Bragg scattering (FIG. 16), the PDF analyzes both Bragg and diffusive scattering and provided details on crystalline structure and the structural deviations from perfect crystallinity such as defects, mismatch or disorder of the materials in atomic scale. The peak of atomic pair in PDF vanished at a distance longer than the longest interatomic distance of the materials (coherent length), which decreased from >50 Å to ˜30 Å as δ increases from 0 to 0.29 (FIG. 13b ). The results demonstrated that crystalline order in Na₆₇ MnO_(x) becomes more confined as δ increased. Namely, Na_(δ)MnOx-H₂O with high concentration of sodium cannot sustain long-range crystallinity and became disordered. Reference to long-range crystallinity is reference to order of greater than 5.0 nm (i.e. the position of the atoms repeat in lattice space in a regular array). Accordingly, in the context of the present disclosure, disordered Na_((δ))MnO_(x)—H₂O may be understood as having crystalline order only up to 3.0 nm and ordered is to be understood as having crystalline order greater than 5 nm. See FIG. 13 b.

FIG. 13c shows that as Na concentration (δ) changed, Na_((δ))MnO_(x)—H₂O (Na_(0.13)MnO_(1.74)-H₂O; Na_(0.25)MnO_(1.84)-H₂O and Na_(0.29)MnO₂—H₂O) showed pure phase Mn₅O₈ (δ=0), mixture of Mn₅O₈ and layered MnO₂ (δ=0.13 and 0.25). When δ reached 0.29, a pure triclinic birnessite structure formed with a chemical formula of Na_(0.29(δ))MnO₂—H₂O, where the Na cations and water molecules occupied the interlayer regions of edge-sharing [MnO₆] octahedra with an interplanar distance of 7.14 Å.

FIG. 17 shows the PDF in the range from 0.8 to 4.3 Å, revealing a phase transition from Mn₅O₈, to a mixture of Mn₅O₈ and MnO₂ birnessite and finally to MnO₂ birnessite as δ increased. The peaks of PDF can be indexed as O—H pair at 0.95 Å (P1) from water (a), Mn—O pairs around 1.9 Å (P2) from the [MnO₆] octahedral unit and 2.2 Å (P3) from Mn atoms in prismatic sites relative to O, Mn—Mn or O—O pair around 2.8 Å (P4), and Mn—O pair around 3.5 Å (P5) from the nearest neighbors of [MnO₆] octahedral units. It is notable that O—H pair (P1) and Mn—O pairs (P2, P3, P5 and P7) showed negative peaks due to negative coherent neutron scattering lengths of H and Mn atoms (−3.74 femtometer and −3.73 femtometer, respectively). The Mn—O pair around 1.9 Å (P2) is attributed to Mn5O8 (b) and layered Na_(0.29)MnO₂ (c), respectively. The Mn—O pair around 2.2 Å (P3) is attributed to Mn(II)-O from Mn₅O₈ phase (d), which decreased relatively to Mn(IV)-O pair at P2 as δ increased, congruent with the decreasing phase fractions of Mn₅O₈. The positive peaks at 2.8 Å (P4) is attributed to Mn—Mn or O—O bonding from adjacent [MnO₆] octahedral units in Mn₅O₈ (e) and Na_(0.29)MnO₂ (f) phases. Therefore, as δ increased the intensity of the PDF peaks at 2.8 Å of each material did not change significantly relatively to Mn(IV)-O pair. Similar trends can be found in Mn—O pair at 3.5 Å (P5) from adjacent [MnO₆] in Mn₅O₈ (g) and Na_(0.29)MnO₂ phases (h). The peaks at ˜4.0 Å (P6 and P7) showed a rather interesting transition from positive to negative direction as δ increased. The positive peak at 3.96 Å (P6) related to O—O pair (i) in Mn₅O₈ either within the same [Mn(IV)O₆] octahedral unit or [Mn(II)-O] units where Mn²⁺ located in the trigonal prismatic site. In contrast, negative peak at 4.0 Å (P7) is attributed to Mn—Na pair at 4.11 Å (h) from the interaction between Na-ion at the interlayer and Mn⁴⁺ from [MnO₆] octahedral unit or Mn—O pair at 3.73 Å (h) from the interaction between H₂O at the interlayer and Mn⁴⁺, both from Na_(0.29)MnO₂ layered phase. The interplay, between negative peaks of Mn—Na and Mn—O_(w) pairs (O_(w) from interlayer H₂O) in Na_(0.29)MnO₂ phase and positive peak of O—O pair from Mn₅O₈ phase at around 4.0 Å, explained the overall peak changed from positive to negative direction when δ increased, again reflecting the phase transition from Mn₅O₈ to Na_(0.29)MnO₂ birnessite driven by the Na-ion insertion during the solid-state annealing.

Based on above analysis, a formation mechanism of Na_(0.29)MnO₂—H₂O birnessite is proposed in FIG. 18, where Mn₃O₄ was converted into Mn₃O₈ through oxidation of [Mn(III)O₆] octahedra of Mn₃O₄ into [Mn(IV)O₆] units, followed by Na-ion driven conversion from Mn₃O₈ to Na_(0.29)MnO₂—H₂O birnessite during thermal annealing in the air. Although Mn₃O₈ and Na_(0.29)MnO₂—H₂O have different crystalline structures, where the former is crystalline monoclinic and the latter is disordered triclinic, both compounds share similar structural characteristics. Mn₃O₈ has a layered structure and consists of sheets of [Mn₃ ⁴⁺O₈]⁴⁻ in the be plane. Each Mn⁴⁺ atom is coordinated by six oxygen atoms and form edge-sharing octahedral unit ([MnO₆]). Half of the Mn⁴⁺ sites in the main octahedral sheets are not fully occupied, above and below these vacant sites are Mn²⁺ sites. Therefore, the negatively charged octahedral sheets are further neutralized and held together by Mn²⁺ atoms located between layers, giving a compositional formula of Mn²⁺ ₂Mn⁴⁺ ₃O₈. Unlike Mn²⁺ ions, Mn' ions have larger radius and thus show trigonal prismatic coordination with oxygen atoms. It is apparent that the [Mn₃ ⁴⁺O₈]⁴⁻ sheets resemble the structure of Na_(0.29)MnO₂—H₂O birnessite comprised of infinite [MnO₆] octahedral layer with intercalated Na cations in between. The transition from Mn₅O₈ to Na_(0.29)MnO₂—H₂O birnessite is an equivalent process of ion-exchange of Mn²⁺ ions in the Mn₂ ²⁺Mn₃ ⁴⁺O₈ with Na⁺ ions in solid state.

Without being limited, it is believed that the Mn²⁺ ions with trigonal prismatic coordination located between the interlayer of Mn₅O₈ had higher mobility than the Mn⁴⁺ ions within octahedral coordination. Accordingly, the insertion of Na-ions into the Mn²⁺ site was kinetically favored, accompanied with the migration of Mn²⁺ ions into the vacant sites in [Mn⁴⁺ ₃O₈]⁴⁻ layers, and finally drove the formation of Na_(0.29)MnO₂. XRD showed that anhydrous Na_(0.29)MnO₂ had interlayer distance of 5.58 Å (FIG. 19), very similar to that of Mn₅O₈ (5.2 Å). Upon water intercalation, the resulting Na_(0.29)MnO₂.H₂O showed an increased interlayer distance of 7.14 Å. Note that the Na-ion driven conversion from Mn₅O₈ to Na_(0.29)MnO₂ disclosed herein contrasts the formation of Li—MnO₂ via the ion-exchange between Ca₂Mn₃O₈ (Ca²⁺ ₂Mn⁴⁺ ₃O₈), isomorphic structure of Mn₅O₈ (Mn²⁺ ₂Mn⁴⁺ ₃O₈), and molten lithium nitrate. In the formation of Li—MnO₂, Li-ions occupied all the available octahedral sites between the [Mn₃ ⁴⁺O₈]⁴⁻ layers rather than the trigonal prismatic sites occupied by Ca²⁺ in the parent Ca₂Mn₃O₈ compound due to much smaller size of Li⁺ compared with Ca²⁺, resulting in the complete conversion to layered LiMnO₂ with R3m or O3 symmetry.

Electrochemical performance of Na_((δ))MnO_(x)—H₂O were tested in a 0.1 M Na ₂50₄ electrolyte in a three-electrode half-cell using cyclic voltametry (CV) measurements between −1.25 V to 1.25 V (vs Ag/AgCl) at scan rates ranging from 5 to 1000 mV s⁻¹ (FIG. 20). FIG. 21a showed the CVs of Na_(0.29)MnO₂—H₂O, where distinct redox peaks can be observed at all the tested scan rates. As the scan rate increased, the anodic peaks shifts to higher potential from 0.78 V to 1.00 V, while the cathodic peaks shifted to lower potential from 0.12 V to 0.17 V. Compared with other Na_(δ)MnO_(x)—H₂O materials (Na_(0.1.3)MnO_(1.74)-H₂O and Na_(0.25)MnO_(1.84)-H₂O), Na_(0.29)MnO₂—H₂O showed the least peak-shifting, indicating it has a more faciliated redox processes requiring lower overpotential for Na-ion transport (FIG. 22). In FIG. 21b , Na_(0.29)MnO₂—H₂O material showed higher specific capacities compared with other materials at all scan rates with a maximum specific capacity 147 mAh g⁻¹ at a scan rate of 5 mV s⁻¹. To further evaluate the Na-ion transport in Na_(δ)MnO_(x)—H₂O materials, the diffusion coefficient was measured using a current pulse relaxation technique.⁷ As shown in FIG. 23, the relative diffusion coefficients (regarding to Mn₅O₈) of Na_(0.13)MnO_(1.74)—H₂O, Na_(0.25)MnO_(1.84)—H₂O, and Na_(0.29)MnO₂—H₂O were 2.4, 6.6 and 38.7, demonstrating the Na-ion intercalation has less energy barrier in the Na_(0.29)MnO₂—H₂O electrode, congruent with CV data showing less overpotential for ionic trasport for Na_(0.29)MnO₂—H₂O.

Long-term energy and power performance of Na_(0.29)MnO₂—H₂O material were tested in symmetric full-cells for 5,000 galvanostatic cycles at a potential window of 2.5 V. Nearly linear votalge-capacity profiles at all the tested current densities pointed out a single-phase solid solution redox reaction (FIGS. 21c , 24). FIG. 21d shows that the electrode capacities of Na_(0.29)MnO₂—H₂O material varied from 83 mAh g⁻¹ to 24 mAh g⁻¹ as the current density increased from 1 to 10 A g⁻¹, corresponded with the discharge time from 160 s (a C-rate of 23) to 4.5 s (a C-rate of 800). Na_(0.29)MnO₂—H₂O material exhibited an excellent cycle stability up to 5000 cycles without obvious capacity loss, as well as nearly 100% coulombic efficiency and high energy efficiency at different current densities (FIG. 21e ). At the low current densities Na_(0.29)MnO₂—H₂O showed a continuous increase capacity upon cycling. Such behaviour has been attributed to the slow building-up of ionic interface during the initial cycling before the electrode reached its best electrochemical condition. FIG. 21f shows that Na_(0.29)MnO₂—H₂O exhibited the specific energy from 26 to 7.5 Wh kg⁻¹ and the specific power from 625 to 6250 W kg⁻¹. These values are higher or comparable with several aqueous or nonaqueous EES devices, including Panasonic (17500) Li-ion battery (data reported in less than 5 cycles), α-MnO₂, δ-MnO₂ or amorphous birnessites, and tunnel-structured Na_(0.44)MnO₂ and O₃ type NaMnO₂.

The limited capacity for aqueous Na-ion found in typical birnessite is attributed to the limited potential window (˜1.2 V) and ineffective redox process. In order to elucidate the origin of high capacity found in Na_(0.29)MnO₂—H₂O birnessite (147 mAh g⁻¹), the roles of disordered nature on increasing the voltage window and therefore inhibiting the gas evolution reaction is considered, as noted below.

To determine whether the structure found in Na_(0.29)MnO₂—H₂O affected the voltage window for aqueous Na-ion storage, CV measurement and Tafel analysis for hydrogen evolution reaction (HER) and oxygen evolution reaction (OER) were conducted using disordered Na_(0.29)MnO₂—H₂O, high-temperature treated Na_(0.29)MnO₂—H₂O made via thermally treating disordered Na_(0.29)MnO₂—H₂O at 500° C., and commercial anhydrous MnO₂ bulk materials. Compared with high-temperature treated Na_(0.29)MnO₂—H₂O and MnO₂, disordered Na_(0.29)MnO₂—H₂O showed much weaker HER current at a potential of up to −1.25 V (equivalent to overpotential of 0.63 V towards HER) and higher Tafel slopes at various scan rates (FIGS. 25a, 25b , 26), suggesting a sluggish kinetics of HER. FIG. 25b showed that although high-temperature treated and disordered Na_(0.29)MnO₂—H₂O were both inactive towards oxygen evolution reaction (OER) even at a potential of 1.25 V (equivalent to overpotential of 0.63 V towards OER), only disordered Na_(0.29)MnO₂—H₂O showed high overpotential towards HER and OER, suggesting that the disordered nature lead to high resistance to gas evolution reactions and therefore a kinetically stable potential window of 2.5 V in an aqueous electrolyte. In contrast, high-temperature treated Na_(0.29)MnO₂—H₂O and commercial MnO₂ showed much inferior capacities compared with disordered Na_(0.29)MnO₂—H₂O (FIG. 25c ), which could also be explained by the parasitic gas evolution reactions especially HER that could deteriorate the electrode and cause capacity loss at prolonged cycles.

WORKING EXAMPLES

Material synthesis. Mn₃O₄ nanoparticles were first synthesized via a solution phase method. In a typical synthesis, MnCl₂.4H₂O (0.7 g, Alfa Aesar, 99% metals basis) was fully dissolved by deionized water (140 mL, 18.2 MΩ; Millipore, Inc.) in a 500 mL flask under vigorous stirring at room temperature. The aqueous solution of NaOH (Alfa Aesar, 99.98% metals basis) with a concentration of 0.123 g mL⁻¹ was injected at a rate of 0.167 mL min⁻¹ for 50 min using an automatic syringe (HSW Inc.). After injection, the mixture continuously reacted for another 30 min till dark brown precipitate was formed. The resulting product was separated by centrifuging and then washed by deionized water and ethanol consecutively. The obtained products (Mn₃O₄ nanoparticles) were finally vacuum-dried.

In the synthesis of sodium-manganese oxides, NaOH (Alfa Aesar, 99.99% metals basis) and 100 mg Mn₃O₄ nanoparticles were ground in mortar with the molar ratios of 0.5, 1, 1.5, 2 and 4, respectively. The resulting mixture of NaOH and Mn₃O₄ was heated in tube furnace (Thermal Scientific, Inc.) in the open air at 270° C. for 6 hours. The obtained solids were thoroughly washed with deionized water to remove the possible NaOH residual and vacuum-dried for overnight. The high-temperature treated Na_(0.29)MnO₂—H₂O material was obtained by thermal treatment of the as-synthesized disordered Na_(0.29)MnO₂ at 500° C. for 2 hours in the open air. The MnO₂ birnessite with low sodium concentration was synthesize via a wet chemistry method. Aqueous MnCl₂ (5 mg mL⁻¹) precursor was injected into 20 mL NaOH solution with a concentration of 5.7 mg mL⁻¹ at the rate of 0.167 mL min⁻¹ for 1 hour, and the obtained product was vacuum-dried after washed by deionized water and ethanol. And then the solids was annealed in the open at 270° C. for 6 hours.

Electrochemical Measurements

Half-cell test: Cyclic voltammetry (CV) measurements of sodium-manganese oxide were conducted using a three-electrode half-cell powered by CHI 660d single channel electrochemical workstation. The three-electrode system contained a glassy carbon rotating disc electrode (Pine Instrument) as the working electrode, platinum wire and silver-silver chloride (Ag/AgCl) electrode as counter and reference electrodes, respectively. The ink material was prepared by grinding mixture of 7 mg sodium-manganese oxide and 3 mg carbon black (Alfa Aesar, >99.9%). The resulting mixture was mixed with deionized water to make an ink solution of 0.5 mg mL⁻¹. The resulting solution was subsequently sonicated until the materials were homogeneously dispersed. In a typical half-cell measurement, 10 μL suspension containing 3.5 μg sodium-manganese oxide and 1.5 μg carbon black was drop-cast onto the glassy carbon disc electrode (0.5 cm in diameter) and vacuum-dried. The CV measurements of electrodes were conducted in a 250 mL flat-bottom flask containing 100 mL argon-purged Na₂SO₄ aqueous electrolyte (0.1 M) at a rotating rate of 500 rpm. The CV data were obtained within an applied potential range from −1.25 V to 1.25 V (vs Ag/AgCl) for 3 cycles, and the third CV cycle was used for the calculation of storage capacity.

Diffusivity Measurements

The diffusivity measurements was tested in a typical half-cell setting as described above, except 40 ug active materials sodium-manganese oxides was loaded on working electrode and 0.25 M Na₂SO₄ was used as electrolyte. A constant negative current pulse of 1 uA was first applied to working electrode and was held for 15 seconds to discharge the electrode from the open circuit potential. After that, the working electrode was relaxed and potential changes were collected for another 1000 seconds.

Full-Cell Test

Symmetric two-electrode full-cells with Na_(0.29)MnO₂—H₂O electrodes were assembled and measured to characterize the energy/power performance and the long cycle stability as well. Electrodes were made by drop casting the slurry containing ˜5 mg Na_(0.29)MnO₂—H₂O and 1.25 mg carbon black as a mass ratio of 4:1 on Toray carbon paper (E-Tek, Inc., 1.5 cm in diameter). The resulting electrodes were weighed with an accurate mass loading of active material after vacuum-dried over-night. Two symmetric electrodes were separated by cellulose-based filter paper (Whatman), and 150 μL Na₂SO₄ aqueous solution (1 M) was used as the electrolyte. The cell stack of electrodes and separator was tightened by stainless plate and compression spring to ensure good electrical contact, and then assembled in the split button-cells (model: EQ-STC, MTI Corp.). Galvanostatic charge and discharge measurements of symmetric full-cells were conducted on the battery analyzer (model: B-TG, Arbin Instruments) within 2.5 V potential window for 5000 cycles at the constant current densities of 1, 2, 5 and 10 A g⁻¹. All the electrochemical calculations are provided in the supporting information.

X-Ray and Neutron Scattering Characterization

X-ray diffraction measurements were conducted at 17-BM-B at the Advanced Photon Source at the Argonne National Laboratory with a wavelength of λ=0.72768 Å. In-situ XRD of electrochemical half-cell measurements were conducted in a home-made cell consisted of thin carbon paper (E-Tek, Inc.) as working electrode, platinum wire and micro Ag/AgCl electrode as counter and reference electrodes, respectively. The Na₂SO₄ aqueous electrolyte (1 M) was used as the electrolyte. The suspension of a mixture of Na_(0.29)MnO₂ and carbon black was drop cast on the thin carbon paper, and then dried naturally in air. The cellulosed based filter paper was used as separator. The cell was then assembled for X-ray measurements. In-situ XRD tests were performed during CV scans from −1.25 V to 1.25 V (vs Ag/AgCl) at the scan rates of 5 mV s⁻¹. GSAS-II software was used to analyze the structural changes during the charge and discharge processes. The total neutron scattering experiment was conducted at the Nanoscale-Ordered Materials Diffractometer (NOMAD) beamline at Spallation Neutron Source at Oak Ridge National Laboratory. The pair distribution function (PDF) analysis was conducted using PDFgui software.

EDS and TEM Characterizations

Energy dispersive X-ray spectroscopy (EDS) was conducted for elemental analysis by an Amray 3300FE field emission SEM with a PGT Imix-PC microanalysis system at University of New Hampshire. Regular transmission electron microscopy (TEM) images were collected on Zeiss/LEO 922 Omega TEM at University of New Hampshire.

The present invention therefore describes the synthesis of Na-rich Na_(0.29)MnO₂.H₂O via a solid-state reaction between Mn₃O₄ and NaOH. The conversion from Mn₃O₄ spinel to monoclinic Mn₅O₈, and to triclinic Na_(0.29)MnO₂—H₂O birnessite driven by the Na intercalation was confirmed by neutron total scattering experiments and PDF analysis. The O—K edge soft X-ray absorption measurements and Tafel analysis for gas evolution reactions suggested that interplay between Na-ion, structural water and Mn valences found in high-temperature treated Na_(0.29)MnO₂might account for its high overpotential toward gas evolution reactions and thereby the kinetically stable potential window of 2.5 V in an aqueous electrolyte. Moreover, electrokinetic analysis and insitu XRD measurements both pointed to a high electron transfer reaction (0.36 and 0.41 electrons) during charging/discharging processes, benefited from the Na-rich structure. The reported promotional effects of the disordered and Na-rich structure on storage capacity of layered birnessite open up a new strategy to design high capacity electrode materials for aqueous energy storage.

TABLE 1 Summaried atomic ratio A/Mn (A is the cation including Na⁺, K⁺) of MnO₂ birnessite synthesized via a solid-state reaction compared with that of a wet chemistry method and those of other works. A/ Analysis Refer- Number Sythesis Method Mn Tool ence 1 Mn₃O₄ + NaOH, in air, 0.29 EDS This work 270° C., 6 hours, solid- state reaction 2 Mn²⁺ + NaOH in solution 0.17 EDS This work with open air 3 Mn²⁺ + KOH in solution 0.15 EDS Ref. 1 with open air 4 Mn²⁺ + K₂S₂O₈ + 0.1  ICP Ref. 2 NaOH in solution 5 Mn²⁺ + KMnO₄ in 0.06 XPS Ref. 3 hydrothermal reaction, (surface 240° C., 3 hours ratio) 6 KMnO₄ + HCl in auto- 0.12(0.01) EDS & Ref. 4 claved reaction, 140° C., ICP-AES 100 min 7 Mn²⁺ + NaOH, air bubble 0.25 ICP-AES Ref. 5 *Note: the averaged A/Mn ratio of MnO₂ made via wet chemistry methods is about 0.14.

REFERENCES

-   1 Yeager, M. et al. Highly Efficient K0.15MnO2 Birnessite Nanosheets     for Stable Pseudocapacitive Cathodes. The Journal of Physical     Chemistry C 116, 20173-20181, doi:10.1021/jp304809r (2012). -   2 Qu, Q. et al. Electrochemical Performance of MnO2 Nanorods in     Neutral Aqueous Electrolytes as a Cathode for Asymmetric     Supercapacitors. The Journal of Physical Chemistry C 113,     14020-14027, doi:10.1021/jp8113094 (2009). -   3 Wang, J., Zhang, G. & Zhang, P. Layered birnessite-type MnO2 with     surface pits for enhanced catalytic formaldehyde oxidation     activity. J. Mater. Chem. A 5, 5719-5725, doi:10.1039/C6TA09793F     (2017). -   4 Zhu, H. T. et al. Birnessite-type MnO2 Nanowalls and Their     Magnetic Properties. The Journal of Physical Chemistry C 112,     17089-17094, doi:10.1021/jp804673n (2008). -   5 Cai, J., Liu, J. & Suib, S. L. Preparative Parameters and     Framework Dopant Effects in the Synthesis of Layer-Structure     Birnessite by Air Oxidation. Chem. Mat. 14, 2071-2077,     doi:10.1021/cm010771h (2002).

TABLE 2 Refined crystal structural parameters of Mn₅O₈ obtained by using the fitting of neutron scattering data with R_(wp) = 5.93%. The x, y, z and mult indicated the atom positions and atom numbers in the unit cell, respectively. Frac and Uiso represents the occupation and isotropic thermal parameters, respectively. The mult shows the atom numbers in the unit cell. All the corresponding values are provided in the table below. Refined crystal structural parameters of Mn₅O₈ Atom Type x y z frac mult Uiso Mn1 Mn + 4 0.000 0.000 0.500 1.000 2 0.014 Mn2 Mn + 4 0.000 0.258 0.000 1.000 4 0.010 Mn3 Mn + 2 0.277 0.000 0.347 1.000 4 0.021 O1 O − 2 0.890 0.227 0.598 1.000 8 0.015 O2 O − 2 0.899 0.000 0.088 1.000 4 0.015 O3 O − 2 0.395 0.000 0.069 1.000 4 0.015 Space group: C 2/m a = 10.397 b = 5.725 c = 4.882 (Å) α = 90 β = 109.816 γ = 90 size: 0.008 μm V = 273.392 (Å³)

TABLE 3 Refined crystal structural parameters of Na_(0.13)MnO_(1.74)—H₂O (Mn₅O₈ and MnO₂) with R_(wp) = 6.04%, showing the phase fraction of Mn₅O₈ and MnO₂ (by mass) is 56% to 44%. Refined crystal structural parameters of Mn₅O₈ Atom Type x y z frac mult Uiso Mn1 Mn + 4 0.000 0.000 0.500 1.000 2 0.005 Mn2 Mn + 4 0.000 0.269 0.000 1.000 4 0.018 Mn3 Mn + 2 0.262 0.000 0.349 1.000 4 0.016 O1 O − 2 0.888 0.229 0.582 1.000 8 0.026 O2 O − 2 0.901 0.000 0.095 1.000 4 0.015 O3 O − 2 0.405 0.000 0.094 1.000 4 0.031 Space group: C 2/m a = 10.432 b = 5.723 c = 4.876 (Å) α = 90 β = 110.012 γ = 90 size: 0.009 μm V = 273.539 (Å³) Refined crystal structural parameters of MnO₂ Atom Type x y z frac mult Uiso Mn1 Mn + 4 0.000 0.000 0.000 1.000 2 0.029 O1 O − 2 0.386 −0.055 0.131 1.000 4 0.050 Na1 Na + 1 0.544 0.365 0.450 0.147 4 0.006 O2 O − 2 0.604 0.333 0.513 0.602 4 0.106 Space group: C − 1 a = 5.058 b = 2.731 c = 7.387 (Å) α = 87.511 β = 104.993 γ = 91.302 size: 0.004 μm V = 98.471 (Å³)

TABLE 4 Refined crystal structural parameters of Na_(0.25)MnO_(1.84)—H₂O (Mn₅O₈ and MnO₂) with R_(wp) = 9.40%, showing the phase fraction of Mn₅O₈ and MnO₂ (by mass) is 36% to 65%. Refined crystal structural parameters of Mn₅O₈ Atom Type x y z frac mult Uiso Mn1 Mn + 4 0.000 0.000 0.500 1.000 2 0.011 Mn2 Mn + 4 0.000 0.281 0.000 1.000 4 0.020 Mn3 Mn + 2 0.248 0.000 0.327 1.000 4 0.006 O1 O − 2 0.888 0.230 0.595 1.000 8 0.016 O2 O − 2 0.903 0.000 0.094 1.000 4 0.037 O3 O − 2 0.428 0.000 0.080 1.000 4 0.038 Space group: C 2/m a = 10.405 b = 5.740 c = 4.876 (Å) α = 90 β = 109.807 γ = 90 size: 0.005 μm V = 273.954 (Å³) Refined crystal structural parameters of MnO₂ Atom Type x y z frac mult Uiso Mn1 Mn + 4 0.000 0.000 0.000 1.000 2 0.052 O1 O − 2 0.391 −0.044 0.138 1.000 4 0.048 Na1 Na + 1 0.515 0.333 0.450 0.137 4 0.007 O2 O − 2 0.611 0.342 0.516 0.448 4 0.112 Space group: C − 1 a = 5.070 b = 2.739 c = 7.363 (Å) α = 86.894 β = 104.782 γ = 90.886 size: 0.011 μm V = 98.730 (Å³)

TABLE 5 Refined crystal structural parameters of Na_(0.29)MnO₂—H₂O (no Mn₅O₈ was observed) with R_(wp) = 13.45%. Refined crystal structural parameters of Na_(0.29)MnO₂ Atom Type x y z frac mult Uiso Mn1 Mn + 4 0.000 0.000 0.000 1.000 2 0.018 O1 O − 2 0.384 −0.038 0.135 1.000 4 0.052 Na1 Na + 1 0.565 0.161 0.450 0.145 4 0.006 O2 O − 2 0.590 0.330 0.514 0.500 4 0.107 Space group: C − 1 a = 5.048 b = 2.755 c = 7.381 (Å) α = 86.479 β = 104.175 γ = 90.402 size: 0.007 μm V = 99.343 (Å³) 

1-17. (canceled)
 18. A method for forming an electrode having a layered metal oxide/conductive polymer comprising: providing a metal oxide; providing a positively charged polymeric ionomer; providing a negatively charged polymeric ionomer; providing water and mixing the metal oxide, positively charged polymeric ionomer and negative charged polymeric ionomer for a period of time of at least 100 hours to form a layered metal oxide with said positively and negatively charged polymeric ionomer which has in-plane chemical bonding with a dissociation energy of 4 to 7 electron volts.
 19. The method of claim 18 wherein said metal oxide is selected from V₂O₅, LiMnO₂, TiO₂, MoO₂, MoO₃, Nb₂O₅ and LiCoO₂.
 20. The method of claim 18 wherein said positively charged polymeric ionomer comprises poly(3,4-ethylene dioxythiophene).
 21. The method of claim 1 wherein said negatively charged polymeric ionomer comprises poly(styrenesulfonate).
 22. The method of claim 18 wherein said metal oxide and conductive polymer are present at a weight ratio of 1:1 to 8:1.
 23. The method of claim 18 wherein said layered metal oxide conductive polymer is present at a thickness in the range of 1 nm to 30 nm.
 24. The method of claim 18 wherein metal oxide comprises V₂O₅ and said conductive polymer comprises poly(3,4-ethylene dioxythiophene) in combination with poly(styrenesulfonate) and indicates a capacity of greater than 75 mAh/g at a scan rate of 10 mV/s.
 25. The method of claim 18 wherein said layered metal oxide/conductive polymer indicates a capacity in the range of 75 mAh/g to 160 mAh/g at a scan rate of 10 mV/s.
 26. The method of claim 18 wherein said metal oxide comprises LiMnO₂ and said conductive polymer comprises poly(3,4-ethylene dioxythiophene) in combination with poly(styrenesulfonate) and indicates a capacity of greater than or equal to 60 mAh/g at a scan rate of 10-20 mV/sec.
 27. The method of claim 26 wherein said layered metal oxide/conductive polymer indicates a capacity in the range of 20 mAh/g to 70 mAh/g at a scan rate of 10-500 mV/sec. 